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DOI: 10.1007/s11837-016-2252-z
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Sintering Behavior and Microstructure Formation of Titanium
Aluminide Alloys Processed by Metal Injection Molding
JULIANO SOYAMA ,1,2,3MICHAEL OEHRING,1THOMAS EBEL,1 KARL ULRICH KAINER,1and FLORIAN PYCZAK1
1.—Institute of Materials Research, Helmholtz-Zentrum Geesthacht, Max-Planck-Straße 1, 21502 Geesthacht, Germany. 2.—Departamento de Engenharia de Materiais (DEMa), Universi-dade Federal de Sa˜o Carlos, Rodovia Washington Luı´s, km 235, Sa˜o Carlos 13565-905, Brazil. 3.—e-mail: [email protected]
The sintering behavior of metal injection molded titanium aluminide alloys, their microstructure formation and resulting mechanical properties were investigated. As reference material, the alloy Ti-45Al-5Nb-0.2B-0.2C at.% (TNB-V5) was selected. Additionally, two other variations with Mo and Mo + Si additions were prepared: 45Al-3Nb-1Mo-0.2B-0.2C at.% and Ti-45Al-3Nb-1Mo-1Si-0.2B-0.2C at.%. The results indicate that the optimum sintering temperature was slightly above the solidus line. With proper sin-tering parameters, very low porosities (<0.5%) and fine microstructures with a colony size<85 lm could be achieved. Considering the sintering temperatures applied, the phase transformations upon cooling could be described as L + b fi b fi a + b fi a fi a + c fi a2+ c, which was in agreement
with the microstructures observed. The effects of Mo and Si were opposite regarding the sintering behavior. Mo addition led to an increase in the opti-mum sintering temperature, whereas Si caused a significant decrease.
INTRODUCTION
Titanium aluminides are lightweight inter-metallics that show excellent creep resistance.1,2 Consequently, their application as high-temperature materials offers great potential for decreasing fuel consumption and simultaneously improving perfor-mance.3 Applications vary from the conservative aircraft industry4to cost-oriented automobile manu-facturers.5,6The fabrication of titanium aluminides is nonetheless challenging considering their brittle-ness and sensitivity to chemical composition (phase balance is considerably influenced by only slight compositional changes).7Therefore, near-net shape powder metallurgy techniques are attractive pro-cessing methods because they offer good control of chemical composition and lead to fine and homoge-nous microstructures, as well as to the avoidance of extensive machining.8
The majority of powder metallurgy techniques require sintering for the final consolidation of the part. The sintering step including the cooling pro-cess is critical because it controls microstructure formation and can be a source of defects and
dimensional distortions. Furthermore, the sintering behavior is heavily dependent on the metallic powders.9 For instance, pre-alloyed powders sinter differently in comparison to a mixture of elemental powders.10,11 Therefore, in any powder metallurgy technique, the characteristics of the powders deter-mine not only the shaping process parameters but also the sintering behavior.
Typical sintering temperatures of titanium alu-minides are very high (about 1500C).12,13
When compared to more traditional materials, for instance, steels, high sintering temperatures neces-sitate the use of powerful and accordingly expen-sive furnaces. Therefore, different approaches have been investigated to decrease the sintering tem-perature in order to reduce costs of both equipment and processing. It is possible to decrease the solidus line, for instance by the addition of Zr,12 or by adding elements that cause transient liquid phase formation during sintering, which decreases residual porosity.14,15 Obviously, by changing the chemical composition, different phase formation and microstructure evolution take place. Since lowering the sintering temperature might decrease 2017 The Minerals, Metals & Materials Society
the high-temperature mechanical strength, it is important to find a compromise between the two properties.
In this work, different alloy compositions were investigated in regard to sintering behavior, microstructure formation, and mechanical proper-ties. TNB-V5 (Ti-45Al-5Nb-0.2B-0.2C at.%) was selected as reference material due to its balanced mechanical properties and excellent creep resis-tance.16 The objective of varying the composition was twofold: firstly, to study the sintering behavior of these compositions that up to now have not been reported, and secondly, to improve the mechanical properties. The elements selected for these purposes were Si and Mo. Si leads to the formation of hard particles (silicides),17while Mo effects solid solution hardening, but, depending on the heat treatment, also precipitation hardening with B2 particles.18 However, since Mo is a strong b phase stabilizer, its content must be cautiously selected in order to avoid extensive b phase formation that might be harmful to the mechanical properties. Moreover, due to its high melting point, additions of Mo bear the risk of increasing the sintering temperature. The specimen preparation was carried out by metal injection molding (MIM) and sintering was conducted at different temperatures, depending on the composi-tion. An indication of the sintering temperatures was obtained by differential scanning calorimetry (DSC). Lastly, the microstructure formation in sintered MIM specimens has been described and the resulting mechanical properties at room tem-perature evaluated.
EXPERIMENTAL PROCEDURE
Ingots with the composition Ti-45Al-5Nb-0.2B-0.2C and Ti-45Al-3Nb-1Mo-0.2B-Ti-45Al-5Nb-0.2B-0.2C (all composi-tions mentioned in the text are in at.%) supplied by GfE, Nu¨rnberg, Germany, were used as raw mate-rial for the preparation of pre-alloyed powders. The ingots were gas-atomized with argon at Helmholtz-Zentrum Geesthacht, Germany. Only the fraction <45 lm was used to prepare the specimens. The oxygen levels of the atomised powders measured by a melt extraction technique were 845 lg/g oxygen for 5Nb-0.2B-0.2C and 708 lg/g for Ti-45Al-3Nb-1Mo-0.2B-0.2C. Three compositions were pre-pared from each of the pre-alloyed powders: Ti-45Al-5Nb-0.2B-0.2C (reference material), Ti-45Al-3Nb-1Mo-0.2B-0.2C and Ti-45Al-1Mo-1Si-0.2B-0.2C, hereafter called TNB-V5, 1Mo and 3Nb-1Mo-1Si, respectively.
The specimens were prepared by MIM using an Arburg machine, Allrounder 320S, with the geometry of tensile specimens according to ISO 2740. The feedstock was prepared under argon atmosphere by mixing the metallic powder and binder system (10 wt.%) in a double Z-Blade mixer (Femix KM, 0.5 K) for 2 h at 120C. The binder system consisted of paraffin waxes, polyethylene-vinyl-acetate and
stearic acid and was applied at a fraction of 10 wt.% (about 32 vol.%). This binder content was selected because it leads to adequate feedstock viscosity for the MIM process and a better surface quality in the sintered parts.19 The addition of Si powder (Alfa Aesar, 99.5% with particle size <45 lm) was con-ducted during feedstock preparation. After MIM, the paraffin waxes were removed chemically in a bath of hexane for 15 h at 45C, and the remaining binder was removed thermally. The thermal debinding and sintering were carried out in a single furnace run using a cold wall furnace from XERION, Germany, of the XVAC series. Sintering was conducted in high vacuum of approximately 10 4 mbar at different temperatures, depending on the composition, for 2 h. The oxygen and carbon levels were measured in the sintered specimens by means of a melt extraction technique (LECO System). The model TC-436AR was used for oxygen, while the CS-444 was used for carbon analysis.
For microstructural investigation, the sintered specimens were mounted in resin followed by grind-ing and polishgrind-ing procedures. An Olympus PMG-3 light microscope together with the image analysis software AnalySIS pro was used for porosity and colony size measurements. The linear intercept method was applied to determine the colony size. Backscattered electron images were recorded using a scanning electron microscope (DSM962; Zeiss). The powders were analysed by DSC under argon for the investigation of the sintering behavior. The DSC experiments were conducted in a Netzsch DSC404 Pegasus calorimeter using a heating rate of 20 K/ min. Density measurements were carried out by the immersion method (ASTM B311) using an analyt-ical scale, Sartorius LA230S.
The tensile tests were conducted using specimens with a strain gauge of 30 mm at room temperature. A servohydraulic test machine model MTS 810 from Materials Testing Systems with a load cell of 100 kN capacity was used. The strain rate applied was 2.38 9 10 5/s. The elongation was measured by an
extensometer attached to the ridges of the speci-men. Three specimens were tensile tested and the results averaged.
RESULTS AND DISCUSSION Sintering Behavior
The sintering behavior was investigated by DSC, which served as a guide for finding appropriate sintering temperatures, and by conducting various sintering cycles at different temperatures. Figure 1
shows the DSC curve for TNB-V5 measured during heating. The several phase transformations that occur up to the melting point are indicated. Two important points were the b-transus that indicates the lowest temperature where the b phase is stable and the onset of the melting peak where the liquid phase becomes stable (solidus line). Consid-ering that diffusion is faster in the b phase,20the
b-Sintering Behavior and Microstructure Formation of Titanium Aluminide Alloys Processed by Metal Injection Molding
transus temperature is an interesting lower bound-ary for sintering, whereas the limiting upper tem-perature is dictated by the melting peak.
Sintering experiments were conducted with the reference material (TNB-V5) at various tempera-tures ranging from the b-transus temperature up to above the solidus line. The results are shown in TableI. Clearly, sintering conducted completely in the solid state led to low porosity (in the order of 1% at 1480C/2 h). However, extremely low porosities (<0.3%) were possible only with sintering temper-atures slightly above the onset of the melting peak in a near-solidus sintering condition. Higher sinter-ing temperatures (>1510C) led to the formation of pronounced liquid phase fractions that caused dimensional distortion. In contrast to the elemental powder metallurgy approach, the use of pre-alloyed powders leads to lower porosities because the exothermic reactions between Ti and Al cause swelling and Kirkendall porosity does not take place.11 The lower density and heterogeneous microstructures21 due to the reactive sintering reactions are the major drawback for powder met-allurgy-processed titanium aluminides, but they can be completely mitigated by using pre-alloyed powders. However, pre-alloyed powders are harder, and thus more difficult to shape by uniaxial press-ing, requiring >1.2 GPa to achieve an accept-able green strength.22 Consequently, MIM offers a good compromise of processability, sintered density and microstructural homogeneity.
Like any sintered material, MIM-processed tita-nium aluminides are believed to undergo all three stages of sintering. Firstly, the powder particles come together forming necks, then the necks grow, and finally the part becomes dense with little residual porosity left. The major densification mech-anisms are dependent on liquid phase and diffusion (grain boundary, surface and volume) rather than evaporation/condensation, since the sintering atmo-sphere applied was a vacuum. Simultaneously with the heating and cooling stages of the sintering
process, the various solid-state transformations described in the binary phase diagram take place, which can critically affect the sintered microstructure.
Even though the precise discrimination between each stage of sintering is difficult, it can be con-cluded that temperatures as high as 1300C are necessary to achieve the intermediate stage in the case of TNB-V5.7 Nevertheless, heavy diffusing elements such as Mo, Zr, W, Ta, etc. might delay neck formation. The analysis of the sintered densi-ties in TableIshows that the porosity changes from 8.9% to 3.3% for specimens sintered at 1460C and 1470C. Typically, a change in the pore structure from smooth interconnected to isolated occurs at porosities in the order of 5%, which characterises the start of final stage sintering. Consequently, at temperatures around 1460C, the final stage sinter-ing is reached. TableIadditionally shows that little change in the porosity of TNB-V5 takes place between 1490C and 1510C. In fact, the mechanical properties of TNB-V5 sintered at 1490C would probably suffice for many load-bearing applications. Nonetheless, for applications in which the shape of the pores are important, e.g., fatigue or thermal fatigue, higher sintering temperatures are neces-sary to achieve rounder pores.12
Based on the findings that near-solidus sintering in the case of TNB-V5 is more efficient to achieve very low porosity, the sintering temperatures for the alloy variations were selected by determining their melting peaks through DSC analysis, as shown in TableII. The addition of Mo decreased the b-transus temperature by about 45C in com-parison to TNB-V5. Moreover, Mo induced an increase in the onset of the melting peak, which in turn is reflected in increasing the optimum sinter-ing temperature. On the other hand, additions of Si led to a significant decrease (35C) in the onset of the melting peak. Independent of the alloy varia-tion, with proper selection of the near-solidus sintering temperature, very low porosities (<0.2%) could be achieved. With additions of Mo and Mo + Si, a pronounced decrease in the colony size also took place. In the case of Mo + Si addition, the colonies were half the size of the reference material. The effects of adding Mo and Si were opposite in regard to optimum sintering temperatures. Mo has a high melting point and is a strong b-stabilizing element. Therefore, the change in the b-transus temperature, as well as the slight increase in the solidus line, was expected. In addition, Mo is a heavy diffusing atom and it thus hampers mass transport and densification during sintering. Con-sidering that near-solidus sintering temperatures were applied, with an increase in the solidus line, an increase in the optimum sintering temperatures naturally followed. In contrast, with Si addition, the onset of the melting peak was considerably shifted to lower temperatures. Therefore, the specimen 3Nb-1Mo-1Si could be sintered in a near-solidus
condition at a much lower temperature. A possible explanation for the decreased sintering temperature in the case of Si addition could be due to the formation of a transient liquid phase before the formation of silicides. However, apart from the shift in the melting peak, no evidence of a different phase transformation at high temperatures could be observed in in the DSC curves. Despite the lower onset of the melting peak due to Mo addition, 3Nb-1Mo-1Si still showed a b-transus temperature sim-ilar to 3Nb-1Mo. Consequently, the shift in the melting peak is likely a thermodynamic effect rather than kinetic and does not involve a transient liquid phase.
Microstructure Formation
The scanning electron microscopy (SEM) micro-graphs of the sintered specimens are shown in Fig.2. All microstructures were fully lamellar and shared some common features such as pores and boride particles. In the case of TNB-V5, a local enrichment of heavy elements (Nb) was also identified (Fig.2a). In 3Nb-1Mo and 3Nb-1Mo-1Si, these regions were composed of Nb and Mo. Additionally, the b phase could also be found at grain boundary triple points in the specimens containing Mo. The phase fraction was, however, small. For the 3Nb-1Mo-1Si, fine particles that are most probably titanium silicides (Ti5Si3), were additionally identified (Fig.2d).
The sintered microstructures of the alloys investi-gated were all fully lamellar. The sintering temper-atures were in the range of the b + L phase field and consequently passing through the a phase field while cooling took place. The lamellar structure forms as
the a + c phase field is crossed, which happens progressively with the c phase precipitating from the a phase. Further cooling leads to the ordering of the hexagonal a phase, changing it to the a2phase.
Considering that the most appropriate sintering temperatures for achieving very low porosity were slightly above the solidus line, the phase transfor-mations during cooling from the sintering tempera-ture can be described as: L + b fi b fi a + b fi a fi a + c fi a2+ c. In the case of Mo addition (b
stabilizer), the phase formation during cooling can be modified to include the presence of a stabilized b phase (bs): L + b fi b fi a + b fi a + bs fi a
+ c + bs fi a2+ c + bs. The major microstructural
differences between TNB-V5 and the alloy variations tested were the presence of the stabilized b phase and of titanium silicides, which were also formed upon cooling from the sintering temperature. Another important phenomenon is the formation of Nb microsegregations (also including Mo for 3Nb-1Mo and 3Nb-1Mo-1Si), which appeared as a network with bright contrast in BSE pictures. These regions arise due to the b fi a phase transition, when the b-stabilizing elements are rejected for the b phase and remain at the former a/b phase boundaries. In the case of 3Nb-1Mo and 3Nb-1Mo-1Si, the enrichment of b-stabilizing ele-ments caused the b phase to be stabilized at room temperature, remaining mostly at the colony boundaries and triple points.
The refinement of colony size observed with Mo and Mo + Si additions can be explained by the effect of more heterogeneous nucleation sites for the a phase and hindering of grain growth in the single a phase field during cooling. Both effects depend on the presence of an additional phase (borides and Ti5Si3) while the transformation b fi a + b fi a
takes place. Therefore, by adding Mo and Mo + Si, it was possible not only to change the optimum sintering temperatures but also to refine the microstructure due to the different phase transfor-mations upon cooling from the sintering temperature.
Mechanical Properties
The tensile properties of the alloys investigated are displayed in TableIII. The alloy 3Nb-1Mo-1Si showed the highest ultimate tensile stress, while
Table I. Sintering behaviour of TNB-V5. The sintering time was 2 h
Sintering temperature (C) Porosity (%) Sintered density (g/cm3) 1460 8.9 ± 0.6 3.81 1470 3.3 ± 0.3 3.93 1480 1.1 ± 0.1 4.06 1490 0.24 ± 0.04 4.12 1500 0.13 ± 0.1 4.13 1510 0.15 ± 0.04 4.12
Table II. b-transus temperatures, the onset of melting peaks, actual sintering temperature applied, residual porosity and colony size of the alloys investigated
Composition a fi a + b (C) Onset of melting peak (C) Sintering temperature (C) Porosity (%) Colony size (lm) TNB-V5 1438 1497 1500 0.13 ± 0.1 82 ± 3 3Nb-1Mo 1393 1502 1510 0.08 ± 0.05 62 ± 6 3Nb-1Mo-1Si 1392 1462 1470 0.13 ± 0.09 42 ± 4
Sintering Behavior and Microstructure Formation of Titanium Aluminide Alloys Processed by Metal Injection Molding
TNB-V5 and 3Nb-1Mo displayed overlapping error bars. TNB-V5 showed slightly more plastic defor-mation. However, the strain at fracture was the same for all alloys.
Figure3 shows three representative uniaxial tensile stress–strain curves. The small improve-ments in the ultimate tensile strength (UTS) with 3Nb-1Mo and 3Nb-1Mo-1Si in comparison to TNB-V5 are indicated. Despite the change in composition, the elastic behavior was quite similar.
Kim et al.23 showed that there is a direct rela-tionship between colony sizes, tensile strength and ductility for titanium aluminides. The Hall–Petch relationship has also been demonstrated to be valid for a wide variety of titanium aluminides.2 Natu-rally, as the colonies get smaller, the more difficult it becomes for the dislocations to move. Nonetheless, the Hall–Petch relationship is questionable in the case of fully lamellar microstructures because the lamellar spacing also plays an important role. Consequently, the higher UTS of 3Nb-1Mo-1Si can only be partly explained by the finer colony size. Since the specimens were cooled at the same rate after sintering, the lamellar spacing is assumed to be constant. Consequently, an additional contribu-tion to the improved tensile properties is probably
from precipitation strengthening, which was the main difference between 3Nb-1Mo and 3Nb-1Mo-1Si. In comparison to duplex microstructures of extruded high Nb containing titanium aluminides that can achieve UTS >1000 MPa and yet consid-erable ductilities of 1–2%,16 the room-temperature
Fig. 2. SEM micrographs of sintered specimens. (a) TNB-V5, (b) Nb-1Mo, (c) 3Nb-1Mo-1Si, (d) 3Nb-1Mo-1Si, detail of b phase at grain boundary triple points and Ti5Si3.
Fig. 3. Representative tensile stress–strain curves at room temper-ature.
tensile properties of MIM-TiAl alloys are rather unsatisfactory. Nonetheless, when compared to the cast + HIP conditions of similar compositions, the mechanical properties are within the same range. For instance, Ti-46Al-8Nb was reported to show UTS of 482 MPa and 0.2% elongation,24 while the same alloy with a refined microstructure due to boron addition, Ti-46Al-8Nb-1B, showed 620 MPa and 0.8%.25 Obviously, the levels of oxygen and
carbon and the microstructures of the cast + HIP materials are different from materials processed by MIM; however, considering that MIM is a near-net shape process, the as-sintered tensile properties are quite acceptable.
The oxygen and carbon levels of the sintered specimens are shown in TableIV. The reference material showed the highest oxygen content but the lowest carbon content. 3Nb-1Mo and 3Nb-1Mo-1Si showed similar values for both oxygen and carbon. The carbon contents are difficult to control in MIM due to the presence of the organic binder required for shaping, thus producing the scatter found among the different alloys. The pick-up of impuri-ties is inevitable in any powder metallurgy process-ing technique. Before the finished consolidated part is achieved, several steps involving powder han-dling are necessary, which usually contribute to increased levels of oxygen, carbon and nitrogen, in addition to the impurities already present in the starting powders. If the solubility limit of intersti-tial elements is exceeded, precipitation of oxides, carbides and nitrides takes place.16 The starting carbon value (0.2 at.%) accounts for 600 lg/g, there-fore it is clear that carbon pick-up occurred during the process. Moreover, the pre-alloyed powders contained approximately 700–800 lg/g of oxygen, consequently their levels practically doubled in the sintered specimens, which could explain the small plasticity of MIM-processed alloys.
Elements such as oxygen and carbon can signif-icantly influence the mechanical properties either due to solid solutions or through the formation of precipitates. Solely by analyzing their levels in TableIV, it is not possible to discriminate their individual contribution to the mechanical proper-ties. However, since their presence usually involves an increase in tensile and creep resistance, it is reasonable to assume that their effect in the alloys studied was not completely detrimental.
CONCLUSION
The sintering behavior and microstructure for-mation of titanium aluminides prepared by metal injection molding were investigated. The most suit-able sintering temperatures were slightly above the solidus line leading to very low porosity (<0.2%). The effects of Mo and Si on the optimum sintering temperature were opposite. While Mo addition increased the optimum sintering temperature, Si addition significantly decreased it, which was a result of a shift in the solidus line. Furthermore, the changes in chemical composition led to the forma-tion of a stabilized b phase and the precipitaforma-tion of fine particles (titanium silicides, Ti5Si3) which
con-tributed to a refinement of colony size and improved the tensile strength at room temperature.
ACKNOWLEDGEMENTS
The authors would like to thank Andreas Dober-nowsky, Dr. Dapeng Zhao, Dirk Matthiessen, Dr. Frank-Peter Schimansky, Gert Wiese, Dr. Jonathan Paul, Prof. Dr. Michael Dahms, Petra Fischer, Uwe Lorenz, Dr. Victor Vitusevych and Wolfgang Lim-berg.
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Table III. Tensile properties of the alloys investigated at room temperature
Alloy UTS (MPa) Plastic strain (%) Strain at fracture (%) TNB-V5 570 ± 1 0.12 ± 0.01 0.5 ± 0.01 3Nb-1Mo 571 ± 17 0.10 ± 0.03 0.5 ± 0.01 3Nb-1Mo-1Si 607 ± 10 0.08 ± 0.02 0.5 ± 0.0
Table IV. Oxygen and carbon levels measured in the alloys investigated after sintering
Alloy Oxygen (lg/g) Carbon (lg/g)
TNB-V5 1780 ± 28 751 ± 87
3Nb-1Mo 1555 ± 42 1192 ± 8 3Nb-1Mo-1Si 1536 ± 53 1237 ± 64
Sintering Behavior and Microstructure Formation of Titanium Aluminide Alloys Processed by Metal Injection Molding
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